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Nanostructured silicon anode and sulfur cathode for lithium rechargeable batteries

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Nanostructured silicon anode and sulfur cathode for lithium rechargeable batteries

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Content NANOSTRUCTURED SILICON ANODE AND SULFUR CATHODE FOR LITHIUM RECHARGEABLE BATTERIES by Jiepeng Rong A Dissertation Presented to the FACULTY OF THE USC GRADUATE SCHOOL UNIVERSITY OF SOUTHERN CALIFORNIA In Partial Fulfillment of the Requirements for the Degree DOCTOR OF PHILOSOPHY (MATERIALS SCIENCE) May 2015 Copyright 2015 Jiepeng Rong All Rights Reserved ii Dedication In dedication to my parents for making me who I am, and my wife for supporting me all the way! iii Acknowledgments My four-and-a-half-year PhD study in University of Southern California was supposed to be hard and stressful as it was meant to be, yet it turned out to be a joyful and memorable journey because of the advisors, colleagues, and friends I met here. I want to express my sincere appreciation to all of them. First I would like to thank my research advisor, Professor Chongwu Zhou, for his generous support financially and intellectually. Professor Zhou gave me lots of freedom to pursue the projects I am interested in. He has also provided insightful discussion, suggestions and guidance for my research, without which this thesis is impossible. I also want to thank the members of my PhD committee and Qualifying Exam committee, Professor Edward Goo, Professor Wei Wu, Professor Andrea Armani, and Professor Stephen Cronin, for their helpful advices for my research and thesis. I want to thank all my colleagues in Professor Zhou’s group during my study here for their help. Among all of them, I want to express my special appreciation to Mingyuan Ge and Xin Fang. We three join the group together, work together on the same projects, and graduate around the same time. It has been my great honor to work with them. They are the most brilliant and hard-working students I have ever seen. Their enthusiasm and love for research is contagious. I will never forget those long, happy nights we spent together in the lab working on the silane CVD. iv I have been so blessed to have the greatest family on earth. I want to thank my parents for their unconditional love and support. I also want to thank my loving, encouraging and patient wife Jingwei for her belief in me. I cannot wait to spend the rest of my life with you. Thank you. Jiepeng Rong Los Angeles, California October 2014 v Table of Contents Dedication ........................................................................................................................... ii Acknowledgments.............................................................................................................. iii List of Figures ................................................................................................................... vii List of Tables ..................................................................................................................... ix Abstract ............................................................................................................................... x Chapter 1. Introduction to Lithium-ion Battery .............................................................. 1 1.1 Background .......................................................................................................... 1 1.2 Working Principles of Lithium-Ion Battery ......................................................... 3 1.3 Cathode Materials ................................................................................................ 4 1.4 Anode Materials ................................................................................................... 7 1.5 Reference .............................................................................................................. 8 Chapter 2. Coaxial Si / Anodic Titanium Oxide / Si Nanotube Arrays for Lithium-ion Battery Anode ................................................................................................................... 12 2.1. Introduction ........................................................................................................ 12 2.2 Experimental Methods ....................................................................................... 14 2. 3 Results and Discussion ....................................................................................... 15 2.3.1 Electrode Fabrication and Characterization ................................................ 15 2.3.2 Electrochemical Measurements .................................................................. 19 2.4 Conclusion .......................................................................................................... 25 2.5 References .......................................................................................................... 26 Chapter 3. Solution Ionic Strength Engineering as a Generic Strategy to Coat Graphene Oxide (GO) on Various Functional Particles and Its Application in Lithium- sulfur (Li-S) Batteries ....................................................................................................... 30 3.1 Introduction ........................................................................................................ 30 3.2 Results and Discussion ....................................................................................... 31 3.3 Experimental Methods ....................................................................................... 45 3.4 Conclusion .......................................................................................................... 46 3.5 Reference ............................................................................................................ 48 vi Chapter 4. Graphene-Oxide-Wrapped Porous Carbon/Sulfur Composite for High- Performance Lithium-Sulfur (Li-S) Battery Cathode ....................................................... 52 4.1 Introduction ........................................................................................................ 52 4.2 Experimental Methods ....................................................................................... 55 4.3 Results and Discussion ....................................................................................... 57 4.4 Conclusion .......................................................................................................... 70 4.5 Reference ............................................................................................................ 71 Chapter 5. Conclusion and Future Work ...................................................................... 74 5.1 Conclusion .......................................................................................................... 74 5.2 Future Work ....................................................................................................... 75 Bibliography ..................................................................................................................... 77 vii List of Figures Figure 1.2.1 Schematic representative of a lithium-ion battery. ....................................... 3 Figure 2.3.1.1 Schematic of fabrication process of coaxial Si-ATO-Si nanotube structure. ..................................................................................................... 15 Figure 2.3.1.2 Characterization of ATO and coaxial Si-ATO-Si nanotubes. (a) SEM of as-prepared ATO nanotubes on a Ti substrate. (b) TEM image of an as- prepared ATO nanotube. (c) Line scan profile over an ATO nanotube (from point A to B in Figure 2.3.1.2c inset). (d) TEM image of the cross-section view of a coaxial Si-ATO-Si nanotube. (e) TEM image of the side view of a coaxial Si-ATO-Si nanotube. (f) Line scan profile over an coaxial Si-ATO- Si nanotube (from point A to B in Figure 2.3.1.2f inset). ........................... 18 Figure 2.3.1.3 Electrochemical performance of a coaxial Si-ATO-Si nanotube anode. (a) Typical cyclic voltammetry curve comparison of the ATO scaffold and coaxial Si-ATO-Si nanotube, showing the inert nature of ATO with Li + between 0.01V-1V v.s. Li/Li + . (b) Galvanostatic charge-discharge voltage profile between 0.01V-1V vs. Li/Li + for the 1 st , 2 nd , and 50 th cycle at 140 mA/g. (c) Charge-discharge specific capacity and Coulombiceffeciencyv.s. cycle number for 50 cycles at different rates ranging from 140 mA/g to 1400 mA/g. Only the weight of Si is considered for specific capacity calculation. (d) Discharge specific capacity and Coulombiceffeciencyv.s. cycle number at 1400 mAh/g with voltage window between 0.01V-1V v.s. Li/Li + , in which condition it takes around 1 hour to charge or discharge the battery. ........................................................................................................ 21 Figure 2.3.1.4 (a,b) SEM image of an ATO nanotube electrode after 10 charge- discharge cycles and rinsing in acetonitrile and 1M HCl consecutively to remove residue electrolyte and SEI. (d,e) SEM image of coaxial Si-ATO-Si nanotube electrode after 10 charge-discharge cycles treated in the same manner. (c,f) EDX line scan profile of Si (blue), Ti (black) and O (red) over (c) one ATO nanotube (from point A to B in Figure c inset), and (f) one coaxial Si-ATO-Si nanotube (from point A to B in Figure f inset). ........... 23 Figure 3.2.1 Digital camera images of (a) GO dispersed in different solutions at the beginning. (b) GO dispersion after 12 hours, and (c) after adding c-S particles to GO dispersion in b. .................................................................. 33 Figure 3.2.2 SEM images of GO dried directly from (a) 1 M HCl solution, and (b) 1 M NH3•H2O solution. Both scale bars correspond to 1 µm. ........................... 35 Figure 3.2.3 SEM characterization of GO coated onto different particles. (a, b) GO coated commercial sulfur (c-S) particles. (c, d) GO coated synthesized sulfur (s-S) particles. (e) GO coated ball-milled silicon particles. (f) GO viii coated commercial carbon black particles. All scale bars correspond to 1 µm, unless otherwise stated. ....................................................................... 36 Figure 3.2.4 Spectroscopic characterizations. (a) Infrared spectra and (b) Raman spectra of GO, c-S, and GO/c-S. The Raman spectra were taken using a 514 nm laser. ............................................................................................................ 39 Figure 3.2.5 Electrochemical measurements of c-S and GO/c-S as Li-S battery cathode materials. (a) CVs of c-S and GO/c-S at 0.1 mV/s in a potential window from 1.9 to 2.6 V vs Li + /Li 0 . (b) Nyquist plot of impedance measurements of c-S and GO/c-S. (c) Specific capacity at different current rates of c-S and GO/c-S. (d) Galvanic charge-discharge performance and Coulombic efficiency of GO/c-S at 1 C (=1 A/g) for 1000 cycles. Specific capacity calculated based on weight of c-S only (dot line), and total weight of GO/c- S are plotted. ............................................................................................... 43 Figure 4.2.1 Schematic of synthesis process of porous C-S, and porous C-S/GO. ......... 55 Figure 4.2.2 Thermal thermal gravimetric analysis (TGA) of porous C-S and porous C- S/GO. Sulfur is 74.7 wt% in porous C-S, and 69.8 wt% in porous C-S/GO. .................................................................................................................... 57 Figure 4.2.3 Electron microscopy characterization of SiO2 spheres, SiO2/sucrose, porous C, porous C-S, and porous C-S/GO. (a-b) SEM image of commercial SiO2 spheres (a), and SiO2/sucrose (b). (c-d) SEM image (c) and TEM image (d) of porous C. (e-f) SEM image (e) and TEM image (f) of porous C-S. (g-h) SEM image (g) and EDX analysis (f) of porous C-S/GO........................... 59 Figure 4.2.4 Electrochemical characterization of bare sulfur, porous C-S, and porous C- S/GO as Li-S cathode materials. (a) Cyclic voltammetry of bare sulfur, porous C-S, and porous C-S/GO at 0.1 mV/s in a potential window from 1.7 to 2.6 V vs. Li + /Li 0 . (b,c,d) Voltage profiles of selected cycles of bare sulfur (b), porous C-S (c), and porous C-S/GO (d). (e) Galvanic charge−discharge performance of bare sulfur, porous C-S, and porous C- S/GO at 1C (=1673 mA/g). ......................................................................... 61 Figure 4.2.5 Electrochemical characterization of bare sulfur, porous C-S, and porous C- S/GO as Li-S cathode materials. (a) Galvanic charge−discharge performance of bare sulfur, porous C-S, and porous C-S/GO at different C rates (1C=1673 mA/g). (b,c,d) Voltage profiles of different C rates of bare sulfur (b), porous C-S (c), and porous C-S/GO (d). ................................... 65 Figure 4.2.6 SEM Characterization of electrodes before (a,c,d) and after 100 charge/discharge cycles (b,d,f). As-fabricated electrodes using bare sulfur (a), porous C-S (c), and porous C-S/GO (e). Electrodes after 100 charge/discharge cycles using bare sulfur (b), porous C-S (d), and porous C-S/GO (f). ................................................................................................. 68 ix List of Tables Table 1.1.1 Comparison of different type of rechargeable batteries……………..……...2 Table 3.2.1. Comparison of ionic strength of solutions used for GO coating in this study……………………………………………………………..………..32 x Abstract Silicon and sulfur are both attractive electrode materials for next-generation rechargeable lithium batteries because of their abundance, high specific capacity, and low cost. This thesis is mainly focused on the research progress I made on silicon-based anode materials and sulfur-based cathode materials for rechargeable lithium batteries with improved energy density, power density, and cycle life. In Chapter 1, the working mechanism of lithium-ion battery, and commercial anode/cathode materials are introduced, which serve as the background information for the following chapters. Then in Chapter 2, the research progress I made in silicon-based anode is presented. To tackle with the problems of silicon anode, such as poor cyclability and early capacity fading due to significant volume change during lithiation and delithiation process, I fabricated a coaxial silicon / anodic titanium oxide / silicon (Si- ATO-Si) nanotube array structure grown on titanium substrate demonstrating excellent electrochemical cyclability. The ATO nanotube scaffold used for Si deposition has many desired features, such as rough surface for enhanced Si adhesion, and direct contact with the Ti substrate working as current collector. More importantly, the ATO scaffold provides a rather unique advantage that Si can be loaded on both the inner and outer surfaces, and an inner pore can be maintained to provide room for Si volume expansion. This coaxial structure shows a capacity above 1500mAh/g after 100 cycles, with less than 0.05% decay per cycle. xi In Chapter 3 and Chapter 4, two kinds of sulfur-based cathode materials are reported. In Chapter 3, a generic and facile method of coating graphene oxide (GO) on sulfur particles is presented. The applications of sulfur/GO core-shell particles as Li-S battery cathode materials are further investigated and the results show that sulfur/GO exhibit significant improvements over bare sulfur particles without coating. Galvanic charge-discharge test using GO/sulfur particles shows a specific capacity of 800 mAh/g is retained after 1000 cycles at 1 A/g current rate if only the mass of sulfur is taken into calculation, and 400 mAh/g if the total mass of sulfur/GO is considered. Most importantly, the capacity decay over 1000 cycles is less than 0.02% per cycle. The coating method developed in this study is facile, robust, and versatile, and is expected to have wide range of applications in improving the properties of particle materials. In Chapter 4, in order to further improve the power density of sulfur-based cathode material, a method of fabricating graphene oxide (GO) wrapped porous carbon/sulfur composite (porous C-S) for high performance lithium-sulfur (Li-S) battery cathode material is reported. A porous C-S composite using conductive porous carbon as framework and sulfur within its channels as filler is synthesized to generate essential electrical contact to the insulating sulfur, thus achieving high specific capacity. Further graphene oxide wrapping over porous C-S is used to mitigate the problem of intermediate polysulphide dissolution into electrolyte during lithiation/delithiation. The electrochemical performance of GO wrapped porous C-S (porous C-S/GO) is investigated, and the results show that porous C-S/GO exhibits significant improvement over bare sulfur and porous C-S without GO wrapping in terms of specific capacity and xii cycling stability, respectively. Galvanic charge-discharge test using porous C-S/GO shows that a specific capacity of 600 mAh/g is retained after 600 cycles at 1 C (=1673 mA/g) current rate if the total mass of porous C-S/GO is considered. Our study proves that graphene oxide can provide a chemically stable interface between active material and binder to keep active materials attached to electrodes during cycling. This could serve as an important guideline for future sulfur-based cathode materials design. 1 Chapter 1. Introduction to Lithium-ion Battery 1.1 Background Availability of cheap and reliable energy is a great challenge for human beings in the 21 st century. Fossil fuels are still playing a major role in the world’s energy supply based on “Basic Research Need for Electrical Energy Storage” published by Department of Energy in 2007. However, fossil fuels are non-renewable natural resource. Besides, the consumption of fossil fuels for electricity generation or transportation will inevitably introduce problems like global warming from carbon dioxide emission and air pollutions. Enormous amount of efforts have been devoted to replace fossil fuels with other green energy sources, such as hydroelectric power, solar, and wind. A well recognized solution to reduce fossil fuel consumption is to generate electricity from green sources that don’t produce carbon dioxide, which includes nuclear, wind, and solar et al. However, most of the green energy sources, such as solar, wind, and hydroelectric power, suffer from their intermittent and uncontrollable nature, which hindered their wide applications. Electrochemical energy storage systems can work together with green energy generation systems (solar, wind) to smooth short term fluctuations in electricity generation and consumption. Besides, electrochemical energy storage technologies are also important to balance the generation and consumption of electricity from green energy sources. For example, electricity generated from solar energy during the daytime need to be stored efficiently to be used at night. So the development of efficient 2 electrochemical energy storage systems is regarded as the key to realize green energy utilization. Type of rechargeable battery Voltage (V) Energy density (Wh/kg) Power density (W/kg) Cycle life Lead-acid 2.1 30-40 180 500-800 Alkaline 1.5 85 50 100-1000 Nickel-cadmium 1.2 40-60 150 1500 Nickel-metal hydride 1.2 30-80 250 500-1000 Lithium-ion 3.7 150-250 1800 500-10000 Table 1.1.1 Comparison of different type of rechargeable batteries. A comparison of different rechargeable battery systems in table 1.1.1 shows that lithium- ion battery can deliver higher voltage, higher energy and power density, and longer cycle life than other popular battery systems.[1] Lithium-ion battery has been growing rapidly as the power source for portable electronics, such as cellular phones and laptop computers since it’s commercialized in 1991 by Sony. However, to be implemented in electric vehicle and make EV practical for efficient and reliable transportation, next- generation lithium-ion battery with improved energy density, power density and cycle life need to be developed. The enhancement of energy density of lithium-ion batteries can be achieved by developing either high capacity cathode and anode materials, or high voltage cathode materials. Besides, the power density and cycle life are also heavily depending 3 on the electrode materials. So the development of new electrode materials with improved electrochemical performance is regarded as the solution to make EV practical. 1.2 Working Principles of Lithium-Ion Battery Figure 1.2.1 Schematic representative of a lithium-ion battery. Figure 1.2.1 shows the schematic representation of a lithium-ion battery. A lithium-ion battery is an electrochemical device that can convert electric energy into electrochemical energy, and vice versa. There are three major components in a lithium-ion battery system: cathode, anode and electrolyte. For commercial lithium-ion batteries, cathode materials are usually transition metal oxides (LiCoO2, LiMn2O4, LiCo1/3Mn1/3Ni1/3O2, LiFePO4, etc.), and anode are mostly made of graphite. Both cathode and anode are intercalation materials, which have unchangeable host matrix with specific sites for lithium ions to be 4 intercalated in. The working mechanism of lithium-ion batteries involves a reversible insertion or extraction of lithium ions into or out of a lithium intercalation compound during changing/discharging. During the charging process, lithium ions generated by cathode travel through electrolyte, and intercalate into anode materials; and lithium ions travel back to cathode during the discharging process. Cathode and anode are separated by a porous separator membrane made of polypropylene/polyethylene soaked in electrolyte. The electrolyte contains lithium salts (i.e. LiPF6) in alkyl organic carbonates (propylene, ethylene, dimethyl carbonate) at different ratios. The separator can prevent electrical contact between cathode and anode. Meanwhile, the pores in separator can provide channels for lithium ions diffusion between cathode and anode. The typical chemical reactions in lithium-ion batteries based on graphite and LiCoO2 are as follows; Anode: C + x Li + + x e - ←→ LixC6 Cathode: LiCoO2 ←→ Li1-xCoO2 + x Li + + x e - 1.3 Cathode Materials In commercial lithium-ion batteries, cathode materials are typically transition metal oxides, which can be oxided to higher valences when lithium is removed.[2] Some of the commonly used or extensively studied cathode materials are discussed as following. 5 The most commonly used cathode material in lithium-ion batteries is LiCoO2.[3] LiCoO2 has the α-NaFeO2 structure, and the cations are in alternating (111) planes in the distorted rock-salt structure. Although LiCoO2 has been a commonly used cathode material for decades, it has problems such as high-cost and toxic to environment and so on. In addition, LiCoO2 can degrade or even fail when overcharged.[4] There are three major reasons for degradation during cycling. First, cobalt could dissolve into electrolyte when LiCoO2 is delithiated during charging.[5] So in the following cycles, there are less LiCoO2 available to host lithium. The second reason is that after lithium intercalation, CoO2 layer would shear from the electrode surface.[6] That can also consume available sites for lithium storage. The third reason is that there could be changes in lattice parameter with different lithium content.[7] Stress and micro-cracks in LiCoO2 can be formed, and eventually lead to mechanical failure. So lots of research has been carried out for alternatives using other transition metals, such as manganese, nickel, and iron. LiNiO2 is a promising candidate to substitute LiCoO2 because of its low cost and high energy density.[8] LiNiO2 also has the α-NaFeO2 structure, which is the same as LiCoO2. LiNiO2 has 20% higher energy density than that of LiCoO2, but is less stable regarding to the electrochemical performance.[9] In LiNiO2, nickel ions tend to occupy lithium sites, which would slow down lithiation and delithiation process. Adding cobalt to LiNiO 2 has been proved to be an efficient way to improve the stability of LiNiO 2. After addition of cobalt, nickel ions tend to occupy sites in cobalt/nickel plane, rather than lithium plane.[10] Another α-NaFeO2 structured cathode material is nickel and cobalt doped 6 LiMnO2. Among different ratios of the three metal ions, LiNi 1/3Mn1/3Co1/3O2 is most commonly used because of its high capacity, good rate capability, and high operating voltage.[11, 12] The improved performance can be attributed to the increased conductivity and structural stability by the addition of cobalt ions.[13, 14] Besides, nickel doping can stabilize the structure and improve the cycling performance.[15] Another promising cathode material is LiMn2O4, which is in spinel structure. In LiMn2O4, the lithium transport path is a 3-D network instead of planes as for the α- NaFeO2 case. LiMn2O4 is also low-cost and environmental friendly comparing to LiCoO2. However, it has a lower specific capacity and serious capacity degradation during charging/discharging. The reason for capacity decay is mainly that manganese ions can dissolve into electrolyte during cycling.[16] To stabilize the structure of LiMn2O4, nickel is commonly added. Nickel is also proven to decrease the lattice parameter and electrical conductivity, and improve the electrochemical performance.[17] Another promising category of cathode materials is LiMPO4 (M=Fe, Co, Mn) with the olivine structure. Among the phosphate-based cathodes, LiFePO4 is the most commonly studied and used material. During delithiation, lithium ions are removed, and Fe 2+ are oxidized to Fe 3+ to maintain charge neutrality.[18-20] Unlike most other cathode materials, LiFePO4 has a very flat discharge voltage profile, which is because there is a miscibility gap between FePO4 and LiFePO4.[21, 22] During delithiation, the two-phase front grows. The formation of the two-phase mixture can deliver a fixed activity, thus an almost constant voltage output.[23-25] LiFePO4 has very low electronic conductivity of 7 10 -9 S cm -1 .[2] In real application, LiFePO4 is either used in nanoparticles or with conductive additives or coating to improve electrical conductivity.[26-28] 1.4 Anode Materials Carbon is the most popular anode material in commercial lithium-ion batteries. The advantages of carbon anode are ease of availability, stability in thermal and chemical environment, good lithium intercalation/de-intercalation reversibility, and low cost.[29- 31] Carbon anode materials can be produced using various raw materials, such as graphite, oil pitch, coal tar, and different resins. These anode materials can fall into three categories (1) graphite, (2) hard carbon, which is non-graphitized glass-like carbon. Hard carbon does not graphitize even if it is treated at high temperature, and (3) soft carbon, which can be converted under high temperature treatment. Usually, graphite and hard carbon are used as anode materials in lithium-ion batteries. At room temperature, after lithiation, graphite converts to LiC6, which is corresponding to a specific capacity of 372 mAh/g. However, there are side-reactions can happen on carbon anode.[32-34] The lithium salt, such as LiPF6, can react with moisture and form HF. HF can cause erosion of active materials and capacity degradation upon cycling. This will result in decomposition of electrolyte and electrode materials. At the same time, a passive layer on carbon anode will form during cycling, which is known as Solid Electrolyte Interface (SEI). Besides, carbon anode has relatively low specific capacity. So, new anode materials with higher specific capacity and more stable cycling performance are in demand. 8 1.5 Reference [1] Scrosati, B. Challenge of portable power. Nature 1995, 373, 557-558. [2] Whittingham, M. S. Lithium Batteries and Cathode Materials. Chemical Reviews 2004, 104, 4271-4302. [3] Antolini, E. LiCoO2: formation, structure, lithium and oxygen nonstoichiometry, electrochemical behaviour and transport properties. Solid State Ionics 2004, 170, 159-171. [4] Belov, D.; Yang, M.-H. Failure mechanism of Li-ion battery at overcharge conditions. J Solid State Electrochem 2008, 12, 885-894. [5] Amatucci, G. G.;Tarascon, J. M.; Klein, L. C. Cobalt dissolution in LiCoO2-based non-aqueous rechargeable batteries. Solid State Ionics 1996, 83, 167-173. [6] Amatucci, G. G.;Tarascon, J. M.; Klein, L. C. CoO2, The End Member of the Li x CoO2 Solid Solution. Journal of The Electrochemical Society 1996, 143, 1114-1123. [7] Ohzuku, T.; Ueda, A. Solid ‐State Redox Reactions of LiCoO2 (R3̅m) for 4 Volt Secondary Lithium Cells. Journal of The Electrochemical Society 1994, 141, 2972-2977. [8] Yamada, A.;Chung, S. 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K.;Sclar, H.;Szmuk Framowitz, Z.;Kovacheva, D.;Saliyski, N.;Gofer, Y.;Sharon, P.;Golik, E.;Markovsky, B.; Aurbach, D. A comparative study of electrodes comprising nanometric and submicron particles of LiNi0.50Mn0.50O2, LiNi0.33Mn0.33Co0.33O2, and LiNi0.40Mn0.40Co0.20O2 layered compounds. Journal of Power Sources 2009, 189, 248-255. [13] Oh, S. W.;Myung, S.-T.;Kang, H. B.; Sun, Y.-K. Effects of Co doping on Li[Ni0.5CoxMn1.5−x]O4 spinel materials for 5V lithium secondary batteries via Co-precipitation. Journal of Power Sources 2009, 189, 752- 756. [14] Wu, H. M.;Tu, J. P.;Yuan, Y. F.;Xiang, J. Y.;Chen, X. T.;Zhao, X. B.; Cao, G. S. Effects of abundant Co doping on the structure and electrochemical characteristics of LiMn1.5Ni0.5−xCoxO4. Journal of Electroanalytical Chemistry 2007, 608, 8-14. [15] Rosciano, F.;Colin, J.-F.;La Mantia, F.;Tran, N.; Novák, P. Electrochemical Stress at High Potential to Investigate Phase Transitions in Li1.1 ( Ni1 / 3Mn1 / 3Co1 / 3 ) 0.9O2. Electrochemical and Solid-State Letters 2009, 12, A140-A144. [16] Yunjian, L.;Xinhai, L.;Huajun, G.;Zhixing, W.;Qiyang, H.;Wenjie, P.; Yong, Y. Electrochemical performance and capacity fading reason of LiMn2O4/graphite batteries stored at room temperature. Journal of Power Sources 2009, 189, 721-725. [17] Fang, T.-T.; Chung, H.-Y. Reassessment of the Electronic-Conduction Behavior above the Verwey-Like Transition of Ni2+- and Al3+-Doped LiMn2O4. Journal of the American Ceramic Society 2008, 91, 342-345. [18] Axmann, P.;Stinner, C.;Wohlfahrt-Mehrens, M.;Mauger, A.;Gendron, F.; Julien, C. M. Nonstoichiometric LiFePO4: Defects and Related Properties. Chemistry of Materials 2009, 21, 1636-1644. [19] Chen, J.;Vacchio, M. J.;Wang, S.;Chernova, N.;Zavalij, P. Y.; Whittingham, M. S. The hydrothermal synthesis and characterization of olivines and related compounds for electrochemical applications. Solid State Ionics 2008, 178, 1676-1693. [20] Maier, J.; Amin, R. Defect Chemistry of LiFePO4. Journal of The Electrochemical Society 2008, 155, A339-A344. [21] Kobayashi, G.;Nishimura, S.-i.;Park, M.-S.;Kanno, R.;Yashima, M.;Ida, T.; Yamada, A. Isolation of Solid Solution Phases in Size-Controlled LixFePO4 at Room Temperature. Advanced Functional Materials 2009, 19, 395-403. 10 [22] Maccario, M.;Croguennec, L.;Desbat, B.;Couzi, M.;Le Cras, F.; Servant, L. Raman and FTIR Spectroscopy Investigations of Carbon-Coated Li x FePO4 Materials. Journal of The Electrochemical Society 2008, 155, A879-A886. [23] Kasavajjula, U. S.;Wang, C.; Arce, P. E. Discharge Model for LiFePO4 Accounting for the Solid Solution Range. Journal of The Electrochemical Society 2008, 155, A866-A874. [24] Allen, J. L.;Jow, T. R.; Wolfenstine, J. Analysis of the FePO4 to LiFePO4 phase transition. J Solid State Electrochem 2008, 12, 1031-1033. [25] Sigle, W.;Amin, R.;Weichert, K.;van Aken, P. A.; Maier, J. Delithiation Study of LiFePO4 Crystals Using Electron Energy-Loss Spectroscopy. Electrochemical and Solid-State Letters 2009, 12, A151-A154. [26] Ojczyk, W.;Marzec, J.;Świerczek, K.;Zając, W.;Molenda, M.;Dziembaj, R.; Molenda, J. Studies of selected synthesis procedures of the conducting LiFePO4-based composite cathode materials for Li-ion batteries. Journal of Power Sources 2007, 173, 700-706. [27] Lin, Y.;Gao, M. X.;Zhu, D.;Liu, Y. F.; Pan, H. G. Effects of carbon coating and iron phosphides on the electrochemical properties of LiFePO4/C. Journal of Power Sources 2008, 184, 444-448. [28] Liu, H.; Tang, D. The low cost synthesis of nanoparticles LiFePO4/C composite for lithium rechargeable batteries. Solid State Ionics 2008, 179, 1897-1901. [29] Marom, R.;Amalraj, S. F.;Leifer, N.;Jacob, D.; Aurbach, D. A review of advanced and practical lithium battery materials. Journal of Materials Chemistry 2011, 21, 9938-9954. [30] Girishkumar, G.;McCloskey, B.;Luntz, A. C.;Swanson, S.; Wilcke, W. Lithium−Air Battery: Promise and Challenges. The Journal of Physical Chemistry Letters 2010, 1, 2193-2203. [31] Scrosati, B.; Garche, J. Lithium batteries: Status, prospects and future. Journal of Power Sources 2010, 195, 2419-2430. [32] Li, H.; Zhou, H. Enhancing the performances of Li-ion batteries by carbon-coating: present and future. Chemical Communications 2012, 48, 1201-1217. [33] Yen, Y.-C.;Chao, S.-C.;Wu, H.-C.; Wu, N.-L. Study on Solid-Electrolyte- Interphase of Si and C-Coated Si Electrodes in Lithium Cells. Journal of The Electrochemical Society 2009, 156, A95-A102. 11 [34] Zhang, H.-L.;Liu, S.-H.;Li, F.;Bai, S.;Liu, C.;Tan, J.; Cheng, H.-M. Electrochemical performance of pyrolytic carbon-coated natural graphite spheres. Carbon 2006, 44, 2212-2218. 12 Chapter 2. Coaxial Si / Anodic Titanium Oxide / Si Nanotube Arrays for Lithium-ion Battery Anode 2.1. Introduction Increasing efforts have been devoted to the development of lithium-ion batteries with higher energy density, higher charging and discharging rate, and longer cycle life to meet the requirement of ever growing portable electronic and next-generation electrical vehicles industry [1-13]. Silicon has drawn particular attention as anode material for lithium-ion batteries primarily because it has the highest known theoretical capacity (3590 mAh/g for Li15Si4, and 4200 mAh/g for Li22Si4), which is nine times higher than that of commercial graphite anodes and other oxide and nitride materials [14]. However, its application is restricted by severe capacity fading caused by pulverization, which is due to large volume expansion and contraction during lithiation and delithiation process (Si + xLi + + xe - ↔LixSi (0≤x≤4.4)). Thus far, it has been evidenced that nanostructured Si could be utilized to obtain improved electrochemical performance over bulk Si and is considered as promising candidates for high performance lithium-ion battery anodes, such as Si nanowires (NWs) [15-17], carbon-Si core-shell NWs [18], carbon-Si nanocomposites [19,20], TiSi2/Si nanonets [21], three-dimensional porous Si [22,23], and sealed Si nanotubes [24]. These studies indicate that the key feature for electrode design is providing enough free space around Si so that large volume expansion could be accommodated. In addition, Si 13 nanostructures grown directly on metallic current collector would enhance both power density and energy density of lithium-ion battery by minimizing internal resistance and the usage of non-active materials, such as carbon black and binder. Here we report a coaxial silicon / anodic titanium oxide / silicon (Si-ATO-Si) nanotube array electrode in which anodic titanium oxide (ATO) nanotubes were rooted on titanium (Ti) current collector as the inert scaffold in the voltage window 0.01V-1V vs Li/Li + , and the outer and inner Si shell worked as active material to store Li + (as shown in Figure 2.3.1.1). This novel coaxial design incorporated all the key features of designing high performance lithium-ion battery anode. ATO scaffold provide a rough surface for improved Si adhesion, and direct contact with the Ti substrate working as current collector. More importantly, Si layers were coated on both outer and inner surface of ATO nanotube array scaffold, leaving abundant space between nanotubes as well as inside each tube, which allowed for better accommodation of volume expansion outward and inward. Experimentally fabricating coaxial Si-ATO-Si nanotube array electrodes and measuring electrochemical performance as lithium-ion battery anode validated the feasibility of our proposal. We have achieved high first discharge (lithiation) capacity of 2824 mAh/g at the current density of 140 mA/g and long-cycling test after that achieved stable cycling performance with capacity over 1500 mAh/g at 1400 mA/g current rate, which meansless than 5% capacity degradation for 100 cycles. Excellent rate capability was also demonstrated. 14 2.2 Experimental Methods ATO nanotube array synthesis The ATO nanotube array was directly formed on a Ti substrate by potentiostatic anodization at 70 V vs a carbon counter electrode in diethylene glycol electrolyte (DEG) containing 1% HF. The anodization was continuously carried out for 19 hours at room temperature. After anodization, the ATO nanotube array on top of Ti substrate was rinsed with ethanol and naturally dried in air. Coaxial Si-ATO-Si nanotube synthesis An amorphous Si layer, functioning as the Li storage media, was deposited on ATO nanotube scaffold by chemical vapor deposition (CVD) of silane (2% SiH 4 balanced in Ar) at 530 o C for 10 min in a quartz tube furnace (1 inch diameter). The total chamber pressure was 100 Torr. Thickness of Si coating could be easily controlled by reaction time. 15 2. 3 Results and Discussion 2.3.1 Electrode Fabrication and Characterization Figure 2.3.1.1 Schematic of fabrication process of coaxial Si-ATO-Si nanotube structure. Coaxial Si-ATO-Si nanotube array electrode was experimentally fabricated by employing a template approach using ATO nanotubes, as shown by the schematic in Figure 2.3.1.1. ATO scaffold was first prepared by anodization of Ti foil [29], and it provided both mechanical support as well as charge transport path for the active amorphous Si layer on both inner and outer surface. In addition, ATO nanotube does not react with lithium in the voltage window 0.01V-1V vs Li/Li + , which has been confirmed by our cyclic voltammetry (CV) measurements (Figure 2.3.1.3a). We successfully got ATO nanotubes with 100 nm to 300 nm in diameter and 1 µm to 10 µm in length by utilizing different reaction time and voltage. Then, an amorphous Si layer, functioning as the Li storage media, was conformally coated on the surface of ATO nanotube scaffold by chemical Ti foil ATO NT Ti foil Coaxial Si-ATO-Si NT Ti foil Anodization Si CVD 16 vapor deposition (CVD). In this way, the weight ratio of the Si layers to the ATO scaffold could be tuned by varying the thickness of Si coating, which could be easily controlled by CVD time. Si layer needs to be thick enough to provide reasonable loading and thinner than the fracture threshold to avoid losing structural integrity during lithiation-delithiation cycling. Different Si layer thicknesses ranging from 20 nm to 80 nm and different ATO nanotube lengths were tested, and no obvious difference in terms of electrochemical performance was observed. The morphology and composition of as-synthesized ATO nanotube array was first characterized by scanning electron microscopy (SEM). Figure 2.3.1.2a shows that vertically aligned ATO nanotube array with uniform diameter could be obtained. The morphology of ATO nanotubes was further characterized by transmission electron microscopy (TEM), as shown in Figure 2.3.1.2b. The ATO nanotubes have an inner radius (Rin) of 150 nm and an outer radius (Rout) of 170 nm. The rough external and internal surface of ATO nanotubes, intrinsically resulted from the anodization process, would be beneficial to enhance the adherence of Si, which was proven to be a key factor to govern the electrochemical performance of lithium-ion batteries [19,30-34]. The element distribution in ATO nanotubes was analyzed by energy dispersive X-ray (EDX) spectroscopy line scan over the cross-section of one nanotube (Figure 2.3.1.2c). The opposite trend of signal intensity between oxygen (O) and Ti over scanning distance is because O signal is solely from ATO, while Ti is from both ATO and Ti substrate beneath, which could explain why Ti exhibits higher signal intensity over the void space in the center. Coaxial Si-ATO-Si nanotubes were fabricated by Si deposition on ATO 17 nanotubes by CVD. Coaxial Si-ATO-Si nanotube fragments were generated by performing sonication on coaxial Si-ATO-Si nanotube array and characterized by TEM, as shown in Figure 2.3.1.2d. Both cross-section view (Figure 2.3.1.2d) and side view (Figure 2.3.1.2e) revealed a coaxial Si-ATO-Si structure with 50 nm uniform conformal Si coating on both inner and outer surface of ATO nanotubes. Si distribution was investigated by EDX (Figure 2.3.1.2f), confirming that Si symmetrically distributed on both inner and outer surface of ATO nanotubes. After 50 nm Si coated on ATO nanotubes, few-nanometers thick Si at the surface will be oxidized to SiOx when exposed to air. Oxygen signals in EDX results could be from both the thin layer of SiO x and ATO nanotube inside. 18 Figure 2.3.1.2 Characterization of ATO and coaxial Si-ATO-Si nanotubes. (a) SEM of as-prepared ATO nanotubes on a Ti substrate. (b) TEM image of an as-prepared ATO nanotube. (c) Line scan profile over an ATO nanotube (from point A to B in Figure 2.3.1.2c inset). (d) TEM image of the cross-section view of a coaxial Si-ATO-Si nanotube. (e) TEM image of the side view of a coaxial Si-ATO-Si nanotube. (f) Line scan profile over an coaxial Si-ATO-Si nanotube (from point A to B in Figure 2.3.1.2f inset). ATO Si (f) (c) (e) (b) (d) (a) 19 2.3.2 Electrochemical Measurements After fabrication of coaxial Si-ATO-Si nanotube array on Ti substrate, we evaluated its electrochemical performance. Figure 2.3.1.3a shows typical cyclic voltammetry (CV) curves of the ATO nanotube array scaffold before and after Si deposition over the voltage window of 0.01V-1V vs Li/Li + at a scan rate of 0.1 mV/s. ATO nanotubes with no Si coating exhibited electrochemical double-layer capacitor behavior with no peak related to reaction with lithium, which confirmed the inactive nature of ATO as a scaffold. Although titanium oxide is a widely studied anode material reacting with lithium as TiO2+ xLi + + xe - ↔LixTiO2 at 1.7 V vs Li/Li + [35], we limited the voltage window between 0.01V-1V vs Li/Li + for both CV and galvanostatic charge/discharge measurements, so TiO2 did not participate in lithiation/delithiation reaction with lithium and hence ATO just functioned as inert scaffold in this study. After Si deposition, the shape of CV curve changed dramatically, and signature peaks of Si-Li alloy/dealloy reactions were observed. The peak at 0.19 V in the cathodic branch (lithiation) corresponds to the conversion of amorphous Si to LixSi. In the anodic branch (delithiation), the two peaks at 0.35 V and 0.52 V are attributed to the delithiation of LixSi back to amorphous Si [36]. Two-electrode 2032 type coin cells were assembled with coaxial Si-ATO-Si nanotube array grown on Ti current collector as working electrode and lithium metal as the counter electrode to investigate its electrochemical performance. No binder or carbon black additives were employed. Teklon ® polymer separator was used in our coin cells. 1.0 M 20 LiPF6 in 1:1 w/w ethylene carbonate/diethyl carbonate was used as electrolyte. Figure 2.3.1.3b shows the voltage profile for the 1 st , 2 nd , and 50 th cycle of galvanostatic charge/discharge measurement at the same current rate of 140 mAh/g. For the first discharge (lithiation) and charge (delithiation), specific capacity reached 3803 mAh/g and 2802 mAh/g respectively, taking only Si mass into calculation. The Coulombic efficiency in the first cycle was 73.7% and approached above 95% at all charge/discharge rates after the first cycle. The limited Coulombic efficiency in the first cycle and the improved performance thereafter could be resulted from the formation of solid electrolyte interphase (SEI) on the electrode surface, which would consume Li + in an irreversible manner. In addition, the silicon oxide (SiOx) formed on Si surface during its exposure to air could also contribute to the restricted Coulombic efficiency in the first cycle. The reaction between SiOx and lithium is partially reversible, and the reversibility depends on the x value following the formula, SiOx +2x Li + ↔ Si + x Li2O [37,38]. Both processes mentioned above mainly happened in the first cycle and were then suppressed or slowed down, thus resulting in the limited Coulombic efficiency in the first cycle and significant improvement afterwards. 21 Figure 2.3.1.3 Electrochemical performance of a coaxial Si-ATO-Si nanotube anode. (a) Typical cyclic voltammetry curve comparison of the ATO scaffold and coaxial Si- ATO-Si nanotube, showing the inert nature of ATO with Li + between 0.01V-1V v.s. Li/Li + . (b) Galvanostatic charge-discharge voltage profile between 0.01V-1V vs. Li/Li + for the 1 st , 2 nd , and 50 th cycle at 140 mA/g. (c) Charge-discharge specific capacity and Coulombiceffeciencyv.s. cycle number for 50 cycles at different rates ranging from 140 mA/g to 1400 mA/g. Only the weight of Si is considered for specific capacity calculation. (d) Discharge specific capacity and Coulombiceffeciencyv.s. cycle number at 1400 mAh/g with voltage window between 0.01V-1V v.s. Li/Li + , in which condition it takes around 1 hour to charge or discharge the battery. 22 As shown in Figure 2.3.1.3c, the coaxial Si-ATO-Si nanotube array electrode exhibited stable cycling performance at different current rates of 140, 280, 700, and 1400 mA/g. The discharge capacities at each current rate were 2717, 2260, 1823, and 1480 mAh/g, respectively. The discharge capacity at 1400mA/g, at which rate it took around 1 hour to fully discharge/charge the battery, was higher than the theoretical capacity of graphite electrode by a factor of four. The long cycle performance of coaxial Si-ATO-Si nanotube array electrode was also explored by continuously charge and discharge at 1400 mA/g for 100 cycles after the first cycle (Figure 2.3.1.3d). The discharge capacity for the second and the 101 st cycle were 1624 and 1548 mAh/g respectively, corresponding to 4.7% degradation for 100 cycles or less than 0.05% decay per cycle, indicating a superior cycling stability of the electrode. The retaining capacity of 1548 mAh/g after 100 charge/discharge cycles was still more than 4 times higher than the theoretical capacity of graphite. Besides the high gravimetric specific capacity achieved, volumetric specific capacity is estimated around 2800 mAh/cm 3 , around 3.5 times of that of graphite (800 mAh/cm 3 ). We estimated the volumetric energy density as follows. According to SEM images, the density of ATO nanotubes is around 200 per 25 µm 2 (5 µm X 5 µm). The length of ATO nanotubes is 10 µm. Density of silicon is 2330 kg/m 3 . The specific capacity of Si anode is 1500 mAh/g as measured. Based on these parameters, we estimate the volumetric energy density to be 2800 mAh/cm 3 , which is 3.5 times of that of graphite (800 mAh/cm 3 ). The electrode also demonstrated favorable Coulombic efficiency, which rapidly recovered to over 98% after the first cycle, and further increased to more than 99% after 10 cycles. 23 Figure 2.3.1.4 (a,b) SEM image of an ATO nanotube electrode after 10 charge-discharge cycles and rinsing in acetonitrile and 1M HCl consecutively to remove residue electrolyte and SEI. (d,e) SEM image of coaxial Si-ATO-Si nanotube electrode after 10 charge- discharge cycles treated in the same manner. (c,f) EDX line scan profile of Si (blue), Ti (black) and O (red) over (c) one ATO nanotube (from point A to B in Figure c inset), and (f) one coaxial Si-ATO-Si nanotube (from point A to B in Figure f inset). (f) (c) 24 The excellent electrochemical performance of the coaxial Si-ATO-Si nanotube array electrode could be attributed to: (1) the ATO nanotube array provided an excellent inactive, mechanically strong scaffold and survived charge/discharge cycling intact. (2) The ATO nanotube scaffold provided direct contact with the Ti substrate working as current collector. This design could enhance both power density and energy density of lithium-ion battery by minimizing internal resistance and the usage of non-active materials. (3) The rough surface and special geometry of ATO nanotubes not only provided an ideal interface for enhanced adhesion between Si and ATO, but also behaved as a superior host of Si with lower stress associated with lithiation. (4) Si could be loaded on both the inner and outer surface of ATO scaffolds, and an inner pore can be maintained to provide room for Si volume expansion. In addition to the justification by battery performance above, these effects could be further confirmed as following. Batteries were disassembled after 10 charge/discharge cycles, and both ATO nanotube electrode and coaxial Si-ATO-Si nanotube array electrode were rinsed with acetonitrile and 1 M HCl consecutively to remove residue electrolyte and SEI. The morphology and composition were characterized with SEM and EDX. Figure 2.3.1.4a,b show vertically aligned ATO nanotube array without Si coating on Ti substrate under different magnification. The wall thickness of ATO nanotubes was comparable to that of as- synthesized ones, providing further evidence to the fact that ATO here functioned as an inactive scaffold. The EDX line scan profile (Figure 2.3.1.4c) of Ti and O over one ATO nanotube shows comparable pattern with that of as-prepared ATO nanotubes as well. As we expected, coaxial Si-ATO-Si nanotube array electrode also showed limited change in terms of morphology and composition after cycling, which agreed with the stable cycling 25 performance we observed. SEM images (Figure 2.3.1.4d,e) and EDX line scan profile of Si (blue), Ti (black), and O (red) (Figure 2.3.1.4f) confirmed the coaxial Si-ATO-Si nanotube array electrode survived charge/discharge cycling intact. 2.4 Conclusion In summary, we successfully fabricated coaxial Si-ATO-Si nanotube array structures and applied this novel structure to lithium-ion battery anode. The coaxial Si-ATO-Si nanotube array demonstrated high specific capacity and excellent cycling performance. After 100 cycles, the capacity still remained above 1500mAh/g and the capacity decay was less than 0.05% per cycle. This excellent cycling stability was due to the unique coaxial structure in which ATO provided a strong inert scaffold, and a rough surface for Si adhesion. This novel structure thus can be a promising candidate for anode material to improve lithium-ion battery performance. 26 2.5 References [1] Scrosati, B. Battery Technology—Challenge of Portable Power. Nature 1995, 373, 557-558. [2] Tarascon, J. 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Solution Ionic Strength Engineering as a Generic Strategy to Coat Graphene Oxide (GO) on Various Functional Particles and Its Application in Lithium-sulfur (Li-S) Batteries 3.1 Introduction Graphene, a monolayer of carbon atoms tightly packed into a two dimensional (2D) honeycomb sp 2 carbon lattice, has drawn significant attention recently because of its high surface area, chemical stability, mechanical strength and flexibility.[1] The unique 2D geometry and excellent properties of graphene and graphene oxide (GO) endow them as one the most commonly used coating materials to form core-shell structured composites, aiming to improve the performance of the core materials for many kinds of applications, such as lithium-ion battery electrode materials[2-4], corrosion inhibitor[5, 6], photocatalysts[7], solar cells[8] ,sensors[9], and drug delivery[10]. Currently, researchers typically utilize surfactants or organic solvents to coat GO onto functional particles, which suffer the shortcomings of high cost and complexity, and so on. This significantly impedes further development of such core-shell structures towards practical applications. Therefore, a more general and robust approach which can achieve highly uniform coating of GO on those particles with arbitrary sizes, geometries, and compositions are highly desired. 31 3.2 Results and Discussion Considering different surface chemistry among particles, such as sulfur, silicon and carbon, it is natural to seek solutions from physical forces, such as electrostatic forces, surface tension, to achieve the above goal. Here, by taking advantage of the high ionic strength in various aqueous ionic solutions, we developed a facile and robust method which is capable of coating GO uniformly on various particles with arbitrary sizes, geometries and compositions. Our method starts with the chemical exfoliation of graphite to prepare GO.[12-21] Although the exact structures of GO are difficult to determine, it is generally believed that GO is rich in epoxides, hydroxyl, keton carbonyls, and carboxylic groups.[22, 23] Among those functional groups anchored to GO, it is believed that the carboxylic groups and hydroxyl groups play key roles in helping GO form stable colloids in water.[24-28] Previous studies on the surface charge (zeta potential) of as prepared GO shows that GO are highly negatively charged when dispersed in water, apparently as a result of ionization of the carboxylic groups and hydroxyl groups that are known to exist on the GO.[19, 29, 30] The studies suggest that the formation of stable GO colloids is attributed to electrostatic repulsion among adjacent GO. It also implies that positive- charged ions in aqueous solution can be attracted onto the surface of negatively charged GO, screen the electrostatic repulsion among GO, and eventually disturb the stable dispersion of GO. 32 DI water Covalent solutions Ionic solutions # 1 2 3 4 5 6 7 8 9 Solute N/A HAc NH 3•H 2O HCl NaOH NaCl NH 4Ac NH 4Cl NaAc Ionic strength 0 M 0.0042 M ~0 M 1 M ~1M ~1M ~1M ~1M ~1M Notes N/A pK a= 4.756[31] PKa= 9.245[32] pK a= -9.3[33] Assuming complete dissociation Table 3.2.1. Comparison of ionic strength of solutions used for GO coating in this study. Inspired by this deliberation, we chose two categories of solutions with significantly different concentrations of ions, i.e., ionic solutions (also known as “electrolyte”) and molecular solutions as dispersing medium to prepare GO suspensions. Ionic strength is widely used to measure the concentrations of ions in solutions, and the ionic strength of used solutions are estimated and compared in Table 3.2.1. The ionic strength is around 1 for ionic solutions, which are more than two orders of magnitude higher than that of molecular solutions used in our experiments. The difference in ionic strength between ionic solutions and molecular solutions indicates that there are sufficient positive and negative charged ions in ionic solutions, while almost no charged ions in molecular solutions. Different ionic strengths result from the nature of different solute compounds. In ionic solutions, solute compounds dissociate into positively charged cations and negatively charged anions easily, when the ionic bonds holding ions together are broken by polar solvents, such as water. By contrast, in molecular solutions, solute compounds stay as neutrally charged molecules after forming solution. 33 Figure 3.2.1 Digital camera images of (a) GO dispersed in different solutions at the beginning. (b) GO dispersion after 12 hours, and (c) after adding c-S particles to GO dispersion in b. 34 We speculate that the availability of charged ions will make a difference in the dispersion of GO in corresponding solutions. GO was first dispersed in different solutions and the results are shown in Figure 3.2.1a. When GO is dispersed in deionized (DI) water (#1), it forms stable colloid for days without precipitation. Similar results were observed in molecular solutions, such as #2 (1 M HAc) and #3 (1 M NH3•H2O). In these solutions, the solutes are in the form of molecules after being dissolved in water. Neutrally charged molecules do not have effect on the electrostatic repulsions among GO, which can still be stable suspension in these solutions. While in ionic solutions, such as electrolyte #4 (1 M HCl), #5 (1 M NaOH), #6 (1 M NaCl), #7 (NH4Ac), #8 (1 M NH4Cl) and #9 (1 M NaAc), the solute compounds are readily dissociated into ions after dissolution. The positive ions will be attracted to the negatively charged GO, screen the electrostatic repulsion between GO, and break the stable dispersion of GO. Precipitation of GO were clearly observed after 12 h rest in all 6 kinds of ionic solutions (Figure 3.2.1b). 35 Figure 3.2.2 SEM images of GO dried directly from (a) 1 M HCl solution, and (b) 1 M NH3 •H2O solution. Both scale bars correspond to 1 µm. GO from both ionic solutions and molecular solutions were dried directly without washing, and characterized using scanning electron microscopy (SEM) (Figure 3.2.2). To minimize the effect of solute compounds on characterization, GO from solution #3 (1 M NH3•H2O) and solution #4 (1 M HCl) are used as examples of each case. As been well known, NH3•H2O and HCl will evaporate away at elevated temperature leaving GO alone. SEM images of GO from ionic solutions showed wrinkled and crumpled morphology. By contrast, GO from molecular solution exhibited a rather flat surface. The different morphologies of dry GO are attributed to their different dispersion morphologies in solutions. Specifically speaking, after GO was dispersed in ionic solutions, the electrostatic repulsive forces among GO was screened by positively charged ions. GO would tend to crumble, and form winkles to minimize its surface energy. The morphology of GO is maintained after direct drying. By contrast, GO is stretched out and form stably dispersed in molecular solutions affected by the negatively charged surface. 36 And neutrally-charged molecules had no effect on it, which could explain the fact GO are spread out on substrate after drying. Figure 3.2.3 SEM characterization of GO coated onto different particles. (a, b) GO coated commercial sulfur (c-S) particles. (c, d) GO coated synthesized sulfur (s-S) particles. (e) GO coated ball-milled silicon particles. (f) GO coated commercial carbon black particles. All scale bars correspond to 1 µm, unless otherwise stated. If GO is the only additive into the ionic solutions, they tend to crumble, form winkle and restacking to minimize their surface energy as shown in Figure 3.2.2. When there are other particles existing in ionic solutions, there will be one extra way for GO to minimize the surface energy, which is coating on the particles next to them, losing the inner side of its surface, and forming a core-shell structure. To verify this assumption, commercial 37 sulfur (c-S) particles in the diameters between 1 µm and 10 µm were used as an example. The particles were added to GO suspensions in solution #1-#9. As expected, ionic solutions and molecular solutions show completely different behaviors. In ionic solutions (#4-9), GO precipitate together with sulfur particles, leaving the upper solution transparent. SEM characterization of the sediments confirms that GO conformally coated on all c-S particles (GO/c-S) (Figure 3.2.3a,b). Weight ratios of GO to c-S in 1:1 (Figure 3.2.3a) and 1:5 (Figure 3.2.3b) were used in the experiments and complete coating has been achieved in both cases. Simply by adjusting the weight ratio of GO to c-S, thickness of GO coating can be tuned. It is noted that, c-S particles in very irregular shapes were also coated with GO conformally as shown in Figure 3.2.3b. The mechanism of the GO coating on c-S particles is that GO will lose electrostatic repulsive force in high concentration ionic solutions, and take hours to precipitate out because of their low density. During this process, if any particles, such as c-S particles, exist in the solution, GO will tend to coat on their surface to minimize their surface energy. However, in molecular solutions, sulfur particles precipitate by themselves because of their high density, while GO still uniformly dispersed in the solutions as a result of electrostatic repulsion among the negatively charged GO. Because the coating process does not involve any chemical reactions, the method can be extended to other particles with different compositions and sizes. To verify this, the same procedures were applied to three other particles, which were synthesized sulfur (s-S) (Diameter ~500 nm), ball milled silicon (Si) particles (Diameter < 500nm), and commercial carbon black (C) particles (Diameter ~100nm). As expected, each of the 38 three kinds of particles precipitated out with GO coating on their outer surface in the ionic solutions, while the particles sediment by themselves in molecular solutions. SEM characterization confirms the complete and uniform wrapping of GO on particles. Figure 3.2.3c and 3.2.3d show the s-S particles coated with GO in low and high magnifications, respectively. It can be seen that sulfur particles aggregated together, forming clusters in the diameters of a few microns, and GO with wrinkles coated on the clusters conformally. Similarly, in Figure 3.2.3e and 3.2.3f, Si aggregates and C aggregates were completely coated with GO, respectively. 39 Figure 3.2.4 Spectroscopic characterizations. (a) Infrared spectra and (b) Raman spectra of GO, c-S, and GO/c-S. The Raman spectra were taken using a 514 nm laser. To study the reaction mechanism and the composition of the products, infrared spectroscopy (IR) and Raman spectroscopy characterization were carried out over c-S, GO and GO/c-S. Figure 3.2.4a shows IR spectra of as-synthesized GO (see Experimental procedures), c-S particles, and GO/c-S. The following functional groups were identified in GO: O-H stretching vibrations (3420 cm -1 ), C=O stretching vibration (1720-1740 cm -1 ), C=C from un-oxidized sp 2 CC bonds (1590-1620 cm -1 ), and C-O vibrations (1250 cm -1 ). The results showed good agreement with literature.[11] The spectrum from c-S showed a rather smooth curve, and no identified signal between 1000 cm -1 and 3700 cm -1 , indicating c-S has no corresponding functional group on its surface. IR spectrum from GO/c-S exhibited exact the same peak positions as that of GO, indicating that all the functional group from GO remain intact after coating, and also confirmed the existence of GO in GO/c-S. This suggests that there were no chemical reactions between GO and S in 40 preparing GO/c-S. The only driving force leads to GO coating on c-S particles should be the tendency of lowing surface energy of GO. To further verify the mechanism of GO coating process, Raman spectroscopy measurements were carried out on GO, c-S, and GO/c-S as well. Raman spectra showed tangential G modes at ~ 1590 cm -1 and disorder-induced D modes at ~1350 cm -1 from both GO and GO/c-S, confirming the existence of GO in both samples. The ID/IG ratios of both GO and GO/c-S were calculated to be around 0.8, indicating that the quality of GO did not change much after coating on sulfur particles. There was no observable peak between 1200 cm-1 to 1700 cm-1 from c-S or silicon wafer substrate used for all Raman characterizations. One of the most important applications of sulfur particles is used as lithium-sulfur (Li-S) battery cathodes. Li-S batteries are promising candidates to power up electric vehicles because of their high theoretical energy density of 2567 W h kg -1 , which is more than 5 times that of lithium-ion batteries based on traditional insertion compound cathodes.[34- 36] Other advantages of Li-S batteries are that elemental sulfur is low cost, low toxic, and abundant. However, the practical application of Li-S batteries is greatly hindered by three major challenges including (1) dissolution of intermediate polysulphide, (2) poor electronic conductivity of S, and (3) large volumetric expansion of S upon lithiation, which result in rapid capacity decay and low Coulombic efficiency.[37-41] Encapsulating sulfur particles with conducting materials, such as graphene oxide,[42, 43] can improve 41 their electronic conductivity and limit polysulphide dissolution simultaneously. Moreover, rich winkles in GO can provide extra space for volume expansion of S upon lithiation and prevent the electrodes from disruption. Thus, we speculate that by using GO/c-S prepared above (GO:c-S=1:1, weight ratio) as cathode material in Li-S batteries, we may solve all the three problems which significantly impede the practical applications of Li-S batteries currently. In order to demonstrate the structural benefits of GO/c-S in improving cathode performance, a series of electrochemical measurements were carried out. As a comparison, pure commercial sulfur powder without GO (c-S) was also tested following the same procedures. Cyclic voltammetry (CV) was used to reveal the electrochemical reaction mechanism of the cathode materials between 1.9 and 2.6 V at a sweep rate of 0.1 mV/s (Figure 3.2.5a). During the first cathodic reduction process, two peaks at 2.24 V and 2.0 V (vs Li + /Li 0 ) were observed. The peak at 2.24 V corresponds to the reduction of sulfur to higher-order polysulfides (Li2Sx, 4<x<8),[44] while the peaks at 2.0 V can be assigned to the reduction of higher-order polysulphides to lower-order polysulphides (Li2Sx, 2≤x≤4).[42, 45] In the following anodic oxidation process, two peaks at approximately 2.3 V and 2.4 V were observed and can be attributed to the conversion of lithium sulfides to polysulphides and sulfur.[2, 46, 47] GO/c-S also has four corresponding peaks, however, at slightly shifted positions. The two anodic peaks were shifted to lower voltage by 0.07 V, while the two cathodic peaks had much smaller shift. The voltage difference between charge and discharge plateaus of GO/c-S was overall much smaller than that of c-S, indicateing that GO coating can help to reduce the 42 polarization and inner resistance of the batteries. Lower polymerization and inner resistance are key factors to achieve long-cycle stability and high power density in batteries and improve their overall performance. To further support the structural benefits of GO/c-S comparing to c-S, electrochemical impedance analysis were conducted on both battery cells from 100 kHz to 10 mHz. The impedance of the cathode in the Li-S batteries depends strongly on the lithium content inside the electrode materials. To maintain uniformity, electrochemical impedance spectroscopy measurements were carried out on the working electrodes at the de-lithiated state after the 1 st cycle. The Nyquist plots obtained are shown in Figure 3.2.5b. The high frequency corresponds to the ohmic serial resistance Rs, which includes both the sheet resistance of the electrode and the resistance of the electrolytes. The semicircle in the middle frequency range indicates the charge transfer resistance Rct, relating to the charge transfer through the electrode/electrolyte interface and the double layer capacity Cdl formed due to the electrostatic charge separation near the electrode/electrolyte interface. Also, the inclined line in the low frequency represents the Warburg impedance W o, which is related to solid-state diffusion of lithium-ions into the electrode material. GO/c-S clearly showed a significantly smaller semicircle than S does, and charge transfer resistance was reduced from 200 Ω to 25 Ω after GO coating on S. In addition, the serial resistance also reduced from 12 Ω to 6.5 Ω after GO coating, indicating a better electrical conductivity of the electrodes. Decreased charge transfer resistance and serial resistance are both favorable to achieve high current rate performance 43 Figure 3.2.5 Electrochemical measurements of c-S and GO/c-S as Li-S battery cathode materials. (a) CVs of c-S and GO/c-S at 0.1 mV/s in a potential window from 1.9 to 2.6 V vs Li + /Li 0 . (b) Nyquist plot of impedance measurements of c-S and GO/c-S. (c) Specific capacity at different current rates of c-S and GO/c-S. (d) Galvanic charge- discharge performance and Coulombic efficiency of GO/c-S at 1 C (=1 A/g) for 1000 cycles. Specific capacity calculated based on weight of c-S only (dot line), and total weight of GO/c-S are plotted. In order to demonstrate the improved electrochemical performance of GO/c-S when it works as a cathode material in Li-S battery, galvanic current measurements were carried 44 out on both GO/c-S and S at different current rates, as shown in Figure 3.2.5c. GO/c-S has slightly lower specific capacity in the first three cycles than that of c-S, owing to the fact that the weight of GO is taken into calculation, but it does not contribute too much capacity. After 10 cycles at 0.1 C rate (1 C=1000 mA/g), specific capacity approaches 600 mAh/g for GO/c-S, and the corresponding Coulombic efficiency is over 99%. In comparison, it is only 350 mAh/g for S under the same test condition. The improvement in cycling stability of GO/c-S was more significant as the current rate increases, as shown in Figure 3.2.5c. GO/c-S showed capacities of 550, 500, 450, 350, and 50 mAh/g at the current rates of 0.2 C, 0.5 C, 1 C, 2 C, and 5 C. By contrast, S only exhibits 200 mAh/g at the current rates of 0.2 C, and neglectable values at all higher current rates tested. Moreover, GO/c-S recovers most of the original capacity when the cycling current is restored to 0.1 C, implying that the structure of GO/c-S electrode remained stable even under high rate cycling. The enhanced cycling stability and high-rate performance can be attributed to the unique structure of conformal coating of the wrinkled GO on c-S. Further galvanic current tests demonstrate that GO/c-S maintains a capacity as high as 400 mAh/g at 1 C over 1000 cycles when the total mass of GO/c-S is taken into calculation (Figure 3.2.5d). Specific capacity based on c-S only is calculated to be around 800 mAh/g after 1000 cycles, which is over 6 times larger than that of commercial metal oxide cathode materials (e.g. LiCoO2=120 mAh/g). It should be noted that the coulombic efficiency was mostly above 99.5% after the first three cycles. The complete coating of GO on c-S prevents sulfur from dissolving into the electrolyte, and results in improved cycling performance. Nevertheless, we found that capacity degradation still occurred to some extent in long cycling test. Therefore, further improvements by combining the 45 method developed here with other strategies such as conductive polymer coating[34, 48], and pore confinement should be interesting to further increase the stability of the electrode. Related work is ongoing in our lab. 3.3 Experimental Methods Graphene oxide (GO) was prepared following the literature.[11] Briefly, a mixture of concentrated H2SO4/H3PO4 (360:40 mL) was added to a mixture of graphite (3.0 g) and KMnO4 (18.0 g). The reaction was kept at 50 o C for 12 h, then cooled to room temperature, and poured into ice (~400 mL) with 3 mL 30% H2O2. The product is centrifuged at 4000 rpm for 1 hour, and the supernatant was decanted away. The GO in the supernatant was washed with water, 30% HCl, and water again using centrifuge. Synthesized sulfur (s-S) particles were synthesized by adding concentrated HCl (0.8 mL, 10 M) to an aqueous solution of NaS2O3 (100 mL, 0.04 M) with the presence of Polyvinylpyrrolidone (PVP, Mw~ 40,000, 0.02 wt%). After reaction for 2 hours, the sulfur particles were washed with ethanol and water, and dispersed into to an aqueous solution. As received commercial sulfur (c-S) powders were grounded with pestle and mortar for 5 minutes before using. 46 GO/c-S was synthesized with different aqueous solutions. GO was first dispersed in certain solution, and then sonicated for 10 minutes. Sulfur particles were added to GO suspension under stirring for 1 hour. The product was centrifuged at 3000 rpm and washed with water and ethanol. GO/c-S was then dried at 60 ℃ in air for 12 hours. Electrochemical measurements. To prepare the working electrodes, various sulfur-based materials were mixed with carbon black (Super P) and polyvinylidene fluoride binder (8:1:1 by weight) in N-methyl-2-pyrrolidinone to form a slurry. The slurry was then coated onto aluminum foil using a doctor blade and dried at 60 ℃ for 12 hours to form the working electrodes. 2032-type coin cells were assembled in an argon-filled glovebox using lithium metal as counter electrode. The electrolyte used was bis(trifluoromethanesulfonyl)imide (1 M) in 1:1 v/v 1,2-dimethoxyethane and 1,3-DOL containing 1 wt% LiNO3. Cyclic voltammetry and galvanostatic cycling were then carried out from 1.9 V- 2.6 V versus Li+/Li0. 3.4 Conclusion A generic method of coating graphene oxide (GO) on particles by engineering the ionic strength of solutions was developed. Uniform coating of GO on various particles with a wide range of sizes, geometries and compositions are achieved. This method provides a facile, robust, and general solution to coat wrinkled GO on different particles in aqueous solution medium. As an example, the application in coating GO on c-S particles as Li-S 47 battery cathode materials was investigated, and we found that the GO/c-S composite material exhibits significant improvements in electrochemical performance over c-S particles without coating. 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A Soft Approach to Encapsulate Sulfur: Polyaniline Nanotubes for Lithium-Sulfur Batteries with Long Cycle Life. Adv. Mater. 2012, 24, 1176-1181. [45] Zhou, G.;Wang, D.-W.;Li, F.;Hou, P.-X.;Yin, L.;Liu, C.;Lu, G. Q.;Gentle, I. R.; Cheng, H.-M. A flexible nanostructured sulphur-carbon nanotube cathode with high rate performance for Li-S batteries. Energ Environ Sci 2012, 5, 8901-8906. [46] Guo, J.;Xu, Y.; Wang, C. Sulfur-Impregnated Disordered Carbon Nanotubes Cathode for Lithium–Sulfur Batteries. Nano Letters 2011, 11, 4288-4294. [47] Zheng, G.;Yang, Y.;Cha, J. J.;Hong, S. S.; Cui, Y. Hollow Carbon Nanofiber- Encapsulated Sulfur Cathodes for High Specific Capacity Rechargeable Lithium Batteries. Nano Letters 2011, 11, 4462-4467. [48] Yang, Y.;Yu, G.;Cha, J. J.;Wu, H.;Vosgueritchian, M.;Yao, Y.;Bao, Z.; Cui, Y. Improving the Performance of Lithium-Sulfur Batteries by Conductive Polymer Coating. ACS Nano 2011, 5, 9187-9193. 52 Chapter 4. Graphene-Oxide-Wrapped Porous Carbon/Sulfur Composite for High-Performance Lithium-Sulfur (Li-S) Battery Cathode 4.1 Introduction Low-cost, high-energy-density and long-cycling-life rechargeable batteries are in high demand for next-generation portable electronics, especially if all-electric vehicles are to be deployed broadly as replacements for gasoline-powered vehicles. Lithium-sulfur (Li-S) batteries are one of the most promising candidates for energy storage devices because of their high theoretical energy density of 2567 W h kg −1 , which is more than 5 times that of lithium-ion batteries based on traditional insertion compound cathodes.[1-7] Other advantages of Li-S batteries are that elemental sulfur is low in cost, low in toxicity, and abundant. However, the practical application of Li-S batteries is greatly hindered by three major challenges, including (1) poor electronic conductivity, and thus low utilization of sulfur, (2) large volumetric expansion of sulfur upon lithiation, and (3) dissolution of intermediate polysulphide into electrolyte, which results in rapid capacity decay and low Coulombic efficiency. A composite using conductive porous carbon as framework and sulfur within its channels as filler was proved to be an effective structure to tackle the first two problems by generating essential electrical contact to the insulating sulfur, thus achieving high specific capacity.[8] Carbon framework infiltrated with sulfur as cathode has achieved high 53 capacity, good rate capability using different carbon materials, such as CMK-3[8], carbon nanotube[9], carbon nanofiber[10, 11], and graphene oxide[2, 12, 13]. However, the third problem, dissolution of intermediate polysulphide, is not fully eliminated because all the conductive carbon materials used above have open-framework structure to allow sulfur to infiltrate into the pores to synthesize carbon-sulfur composites, which would also provide channels for polysulphide to dissolve into electrolyte during charge/discharge. This intrinsic problem prohibits itself from fully solving all three problems using conductive porous carbon only. Graphene, a monolayer of carbon atoms tightly packed into a two-dimensional (2-D) honeycomb sp 2 carbon lattice, has drawn significant attention recently because of its high surface area, chemical stability, mechanical strength, and flexibility.[14] The unique 2D geometry and excellent properties of graphene and graphene oxide (GO) endow them as one of the most commonly used coating materials to form core-shell structured composites, aiming to improve the performance of battery electrode materials[12, 15, 16] We recently reported a facile and robust method that is capable of coating GO uniformly on various particles with arbitrary sizes, geometries, and compositions, by engineering the ionic solution strength in various aqueous solutions.[17] In particular, we produced sulfur/GO core-shell particles as Li-S battery cathode materials showing superior specific capacity of 400 mAh/g after 1000 cycles at 1 A/g current rate if the total mass of sulfur/GO is considered. Encapsulating sulfur particles with graphene oxide can limit polysulphide dissolution, and may be the key to mitigate polysulphide dissolution and may reveal itself in improved cycling stability. 54 In this paper, we combined the two strategies of using porous C as conductive framework and GO wrapping, and designed porous C-S/GO for high-rate, and long-cycling Li-S cathode material. Porous C-S/GO composite uses conductive porous carbon as framework and sulfur within its channels as filler to generate essential electrical contact to the insulating sulfur, thus achieving high specific capacity. GO wrapping in porous C- S/GO is used to mitigate the problem of intermediate polysulphide dissolution into electrolyte during lithiation/delithiation. Porous C-S/GO composite achieved 600 mAh/g after 600 cycles at 1 C (1C=1672 mAh/g), when the total weight of porous carbon, sulfur, and GO are taken into calculation. And there was only 0.04% capacity decay per cycle between the 2 nd cycle and the 600 th cycle, which is comparable to the best results in literature.[2, 18] Considering we used commercial sulfur powder as sulfur source, instead of Na2S [2] or Na2S2O3 [18], our method showed advantages in cost and simplicity. 55 4.2 Experimental Methods Figure 4.2.1 Schematic of synthesis process of porous C-S, and porous C-S/GO. The synthesis process is shown in Figure 4.2.1. We used commercial silicon dioxide (SiO2) as template. 1 g of sucrose was added to 10 ml SiO2 colloidal (SNOWTEX-ZL ® , 40 wt% SiO2 in water). And the mixture was transferred to a Teflon autoclave, and went through hydrothermal treatment at 180 o C for 2 hours to get sucrose-coated SiO2 particles (SiO2/sucrose). The product was taken out of autoclave and dried in air at 90 o C and then annealed in argon at 500 o C for 2 hours to carbonize sucrose and get carbon-coated SiO2 (SiO2/carbon). After annealing, SiO2/carbon was immersed into 20% HF solution overnight to etch away SiO2 template, and get porous carbon (porous C). Porous C was then washed with deionized (DI) water, ethanol, and DI water again to fully remove HF. To infiltrate sulfur into the pores of porous C, porous C and commercial sulfur power were mixed at 1:3 weight ratio, and hand-grounded together. The mixture was then 56 transferred to a Teflon autoclave and treated at 160 o C for 2 hours. Sulfur was melted, and absorbed into the pores in porous C. The product (porous C-S) was taken out of Teflon autoclave, and analyzed with thermal gravimetric analysis (TGA) to determine the sulfur content to be 74.7% (Figure 4.2.2). GO was prepared following a reported method.[19] Briefly, a mixture of concentrated sulfuric acid (360 ml) and phosphoric acid (40 ml) was added to a mixture of graphite (3.0 g) and potassium permanganate (18.0 g). The reaction was kept at 50 o C for 12 hours, then cooled to room temperature, and poured into ice (~ 400 ml) with 3 ml 30% hydrogen peroxide. The product was centrifuged at 4000 rpm for 1 hour, and the supernatant was decanted away. The GO in the supernatant was washed with DI water, 30% hydrochloric acid, and DI water again using centrifuge. GO coating on porous C/S was carried out following our published work.[17] Briefly, as-synthesized GO (8 mg) and porous C-S (100 mg) were dispersed in 1 M HCl under mild stirring. The electrostatic repulsive force between negatively-charged GO sheets[20] keeping GO dispersed were screened by dissociated ions in HCl solution. GO tends to coat on porous C/S particles to minimize surface energy. The sediments, named as “porous C-S/GO”, were collected using vacuum filtration, washed with DI water, and dried. The weight percentage of sulfur in porous C-S/GO was determined with TGA to be 69.8% (Figure 4.2.2). 57 Figure 4.2.2 Thermal thermal gravimetric analysis (TGA) of porous C-S and porous C-S/GO. Sulfur is 74.7 wt% in porous C-S, and 69.8 wt% in porous C-S/GO. 4.3 Results and Discussion The products at each synthesis step were characterized by scanning electron microscopy (SEM) or transmission electron microscopy (TEM). Figure 4.2.3a, b show the SEM picture of commercial SiO2 spheres and SiO2/sucrose, respectively. Both of them show spherical geometry, with diameters of around 125 nm. Because of the relatively wide distribution of SiO2 diameter, it is difficult to tell the size change before and after sucrose coating. After pyrolysis in argon, sucrose was turned into carbon, which proved the existence of sucrose coating on SiO2 spheres. SEM (Figure 4.2.3c) and TEM (Figure 4.2.3d) images of porous C reveal it is made of interconnected carbon shells with openings in the outermost layer of carbon shells. The openings are due to rupture of the 58 thin carbon shells during processing. The wall thickness of carbon shell is about 5 nm, with 125 nm diameter of void, which is consistent with the diameter of SiO2 template. The thin wall of carbon shell ensures the low mass contribution of carbon, while still provides enough electric conduction paths. SEM (Figure 4.2.3e) and TEM (Figure 4.2.3f) images of porous C-S reveal sulfur is filled up into the pores of porous C, as revealed by the disappearance of voids. The difference in contrast between porous C (Figure 4.2.3d) and porous C-S (Figure 4.2.3f) indicates sulfur is successfully infiltrated into porous C. After being coated with GO, porous C-S/GO was again characterized with SEM (Figure 4.2.3g), and energy-dispersive X-ray spectroscopy (EDX) (Figure 4.2.3h). SEM (Figure 4.2.3g) shows GO is conformally coated on porous C/S. Complete coating of GO on porous C/S can significantly slow down the dissolution of polysulphide into electrolyte, while the spacing between GO layers provides channels for lithium ion transport. EDX spectrum (Figure 4.2.3h) collected with an accelerating voltage of 10 keV over porous C- S/GO sample shows dominant sulfur peak, and much smaller carbon and oxygen peaks. The results confirmed the existence and relative high percentage of sulfur, which is in agreement with TGA results (70 wt% sulfur). The oxygen peak is attributed to functional groups in GO. 59 Figure 4.2.3 Electron microscopy characterization of SiO2 spheres, SiO2/sucrose, porous C, porous C-S, and porous C-S/GO. (a-b) SEM image of commercial SiO2 spheres (a), and SiO2/sucrose (b). (c-d) SEM image (c) and TEM image (d) of porous C. (e-f) SEM image (e) and TEM image (f) of porous C-S. (g-h) SEM image (g) and EDX analysis (f) of porous C-S/GO. 60 Porous C-S/GO was then assembled as Li-S battery cathode for electrochemical measurements. To prepare cathodes, porous C-S/GO was mixed with carbon black (super P ® ) and polyvinylidene fluoride binder (8:1:1 by weight) in N-methyl-2-pyrrolidone to form a slurry. The slurry was then coated onto aluminum foil using a doctor blade and dried at 85 o C for 24 hours to form working electrodes. 2032-type coin cells were assembled in an argon-filled glovebox using lithium metal as counter electrode. The electrolyte used was 1 M lithium bis(trifluoromethylsulfonyl)imide in 1:1 v/v 1,2- dimethoxyethane and 1,3-dioxolane containing 1 wt% lithium nitrate. Bare sulfur and porous C-S without GO coating were also tested as Li-S battery cathode materials following the same procedure to fabricate electrodes as control samples. Electrochemical measurements were also carried out on bare sulfur, porous C-S, and porous C-S/GO. 61 Figure 4.2.4 Electrochemical characterization of bare sulfur, porous C-S, and porous C-S/GO as Li-S cathode materials. (a) Cyclic voltammetry of bare sulfur, porous C-S, and porous C-S/GO at 0.1 mV/s in a potential window from 1.7 to 2.6 V vs. Li + /Li 0 . (b,c,d) Voltage profiles of selected cycles of bare sulfur (b), porous C-S (c), and porous C-S/GO (d). (e) Galvanic charge −discharge performance of bare sulfur, porous C- S, and porous C-S/GO at 1C (=1673 mA/g). 62 Cyclic Voltammetry (CV) was used to reveal the electrochemical reaction mechanism of the cathode materials (Figure 4.2.4a). The voltage window and scan rate used were 1.7 V – 2.6 V versus Li + /Li 0 , and 0.1 mV/s, respectively. During the first cathodic reduction process of bare sulfur, two peaks at 2.24 and 1.92 V (vs Li + /Li 0 ) were observed. The peak at 2.24 V corresponds to the reduction of sulfur to higher-order polysulfides (Li2Sx, 4 < x < 8),[21] while the peak at 1.92 V can be assigned to the reduction of higher-order polysulphides to lower-order polysulphides (Li2Sx, 2 ≤ x ≤4).[13, 22] In the following anodic oxidation process, two peaks at approximately 2.4 and 2.56 V were observed and can be attributed to the conversion of lithium sulfides to polysulphides and sulfur.[10, 12, 23] In comparison to the peaks of bare sulfur, porous C-S and porous C-S/GO have two cathodic peaks shifted from 1.92 V and 2.24 V to higher voltage at 1.95 V and 2.28 V, respectively. Also, the anodic peak at 2.56 V for bare sulfur shifted to 2.51 V for porous C-S, and 2.47 V for porous C-S/GO. The shifting of peaks after incorporating porous C and GO reduced the voltage difference between anodic and cathodic peaks, which indicates both porous C and GO can help to reduce the polarization and inner resistance of the batteries, which are key to achieve good rate performance. It is noted that one peak at 2.4 V in anodic branch remained at the same position for both sulfur and porous C- S/GO sample, and was suppressed in porous C/S sample. This peak cannot be explained by the theory above, and the mechanism deserves further study. Besides, broad cathodic and anodic peaks were observed for sulfur. While both porous C-S and porous C-S/GO electrodes exhibited much narrower peaks. Narrow CV peaks imply that the active materials are confined in a highly ion/charge accessible environment, [24-26] which is again favorable for high rate performance. 63 The superior structure of porous C-S/GO was first proved by long cycling test, using bare sulfur and porous C-S as control samples. All three samples were continuously charge- discharged at 1 C (=1673 mA/g) current rate. The voltage profiles of the 1 st , 2 nd , 10 th and 100 th cycle of bare sulfur (Figure 4.2.4b), porous C-S (Figure 4.2.4c), and porous C-S/GO (Figure 4.2.4d) were plotted, respectively. Porous C-S exhibited significant improvement in terms of specific capacity at each plotted cycle over that of bare sulfur. The specific capacity in the 1 st , 2 nd , 10 th , and 100 th discharge were increased from 571, 533, 230, and 52 mAh/g for bare sulfur to 1095, 916, 615 and 289 mAh/g for porous C-S, respectively. Enhancement of specific capacity was resulted from better utilization of insulating sulfur by introducing porous C as an electrical conductive framework. Although porous C-S could deliver satisfying specific capacity in the first cycle, it experienced significant capacity degradation in the following cycles. In the 100 th discharge, porous C-S only kept 26% of the discharge specific capacity of the 1 st cycle, in comparison to 9% retention over 100 cycles for bare sulfur. That can be explained by the fact that porous C is an open framework with openings in the outermost layer of carbon shells, which exposes sulfur to electrolyte after assembling battery cells. Exposed sulfur can become polysulphide, and dissolve into electrolyte. After further wrapping porous C-S with GO, the specific capacity retention improved to 68% after 100 cycles continuous charge-discharge, and 57% after 600 cycles. The improved cycling stability proves GO wrapping is an effective way to protect sulfur cathode from dissolution in the form of polysulphides during charge- discharge. It is noted that coating 7.4% wt. GO onto porous C-S lowered down the 1 st discharge capacity by about 9%, because GO contributed to total mass but did not participate in electrochemical reaction between 1.7-2.6 V vs. Li/Li + . After 100 cycles, 64 porous C-S/GO delivered more than doubled specific capacity than porous C-S did, because of better capacity retention of porous C-S/GO. The results proved the importance of GO wrapping in fabricating long cycling sulfur cathodes in Li-S batteries. A direct comparison of cycling performance of sulfur, porous C-S, and porous C-S/GO are plotted in Figure 4.2.4e. Porous C-S exhibited enhanced specific capacity in comparison to sulfur, but still with some degree of degradation. Porous C-S/GO maintained the enhanced specific capacity of porous C-S, and further improved the stability of capacity. The results again proved porous C framework and GO wrapping are both proved to be indispensable to fabricate high specific capacity and long-lasting sulfur cathodes for Li-S batteries. 65 Figure 4.2.5 Electrochemical characterization of bare sulfur, porous C-S, and porous C-S/GO as Li-S cathode materials. (a) Galvanic charge −discharge performance of bare sulfur, porous C-S, and porous C-S/GO at different C rates (1C=1673 mA/g). (b,c,d) Voltage profiles of different C rates of bare sulfur (b), porous C-S (c), and porous C-S/GO (d). In order to further examine the effectiveness of the porous C framework and GO wrapping in improving the electrochemical performance of sulfur-based cathodes, the rate capability of bare sulfur, porous C-S, and porous C-S/GO electrodes were 66 investigated and compared in Figure 4.2.5a. Bare sulfur electrodes can deliver about 800 mAh/g in the first cycle at 0.1C current rate, but degraded to less than 400 mAh/g after 10 cycles. Besides, bare sulfur can hardly deliver any meaningful capacity at current rates of 0.5C, 1C, 2C and 5C. Again, the results confirmed bare sulfur has low capacity, low sulfur utilization, poor retention during cycling, and poor rate capability. After infiltrating sulfur into porous C framework, porous C-S improved the specific capacity, sulfur utilization, and rate capability significantly. The first cycle specific capacity increased to around 1300 mAh/g at 0.1C. Even at high current rate, like 0.5C, 1C, 2C and 5C, the specific capacity also increased from almost zero for bare sulfur to 670, 524, 438, and 120 mAh/g, respectively. The only remaining problem for porous C-S was the poor capacity retention during cycling, especially at low current rate. At 0.1C, porous C-S and bare sulfur both showed degradation in the first 10 cycles (Figure 4.2.5a). These results also proved the openings in porous C can provide channel for polysulphide to dissolve away, and cause capacity degradation. To overcome the poor capacity retention problem, porous C-S particles were further wrapped with GO. The electrochemical test at different current rates for porous C-S/GO exhibited significant improvement in capacity retention especially at low current rate, such as 0.1C (Figure 4.2.5a). More importantly, because GO was used, specific capacity of porous C-S/GO at 0.1 C was lower than that of porous C-S in the 1 st discharge. However, after 10 cycles of charge/discharge, porous C-S/GO showed higher specific capacity than porous C-S because of better capacity retention. The voltage profiles at different current rates of bare sulfur (Figure 4.2.5b), porous C-S (Figure 4.2.5c), and porous C-S/GO (Figure 4.2.5d) are plotted, respectively. For bare sulfur in Figure 4.2.5b, discharge plateau can be observed only at 0.1 C rate. At 0.5C or 67 higher current rates, the voltage profiles are not observable. For both porous C-S (Figure 4.2.5c) and porous C-S/GO (Figure 4.2.5d), there are obvious plateaus even at 5C. Besides, both porous C-S and porous C-S/GO showed much smaller voltage difference between charge and discharge plateaus than that of bare sulfur, indicating reduced polarization and inner resistance of the batteries, which is important to achieve good rate capability. 68 Figure 4.2.6 SEM Characterization of electrodes before (a,c,d) and after 100 charge/discharge cycles (b,d,f). As-fabricated electrodes using bare sulfur (a), porous C- S (c), and porous C-S/GO (e). Electrodes after 100 charge/discharge cycles using bare sulfur (b), porous C-S (d), and porous C-S/GO (f). 69 To better understand the working mechanism of porous C framework and GO wrapping, bare sulfur electrodes (Figure 4.2.6a,b), porous C-S electrodes (Figure 4.2.6c,d), and porous C-S/GO electrodes (Figure 4.2.6e,f) were examined with SEM before (Figure 4.2.6a,c,e) and after 100 charge/discharge cycles (Figure 4.2.6b,d,f). The as-prepared bare sulfur electrode (Figure 4.2.6a) showed rough surface. The particles with diameters of micrometers are sulfur, and the much smaller particles with diameters below 100 mm, and uniformly coated on the outer surface were carbon black. After 100 charge/discharge cycles, bare sulfur electrode (Figure 4.2.6b) showed that a lot of sulfur particles were missing, leaving pores around 10 µm in diameter in the electrodes. Although sulfur particles were anchored onto as-fabricated electrodes by the binder, during cycling, sulfur at the sulfur-binder interface may become polysulphide, and get dissolved into electrolyte. Once the sulfur-binder interface got destructed, sulfur particles may get detached from electrodes. Similar phenomenon was observed for porous C-S electrodes by comparing SEM images before (Figure 4.2.6c) and after cycling (Figure 4.2.6d). Some porous C-S particles were detached from electrodes after 100 cycles for the same reason. Porous C is an open framework with openings on the outermost layer, and sulfur had direct contact with binder at the opening locations. Dissolution of polysulphide will disrupt the contact between sulfur and binder, and cause porous C-S particles to detach from the electrode. In the porous C-S/GO electrodes, porous C-S particles were fully wrapped with GO. The binder would have direct contact with GO, instead of sulfur. In contrast, GO can provide a chemically stable interface between the active material and binder, which can be confirmed by SEM images of porous C-S/GO before (Figure 4.2.6e) and after cycling (Figure 4.2.6f). Although the electrodes changed from a continuous film before cycling 70 (Figure 4.2.6e) to domains with gaps between them after 100 continuous lithiation/delithiation cycles (Figure 4.2.6f), each domain maintained its strong adhesion onto the underlaid current collector. The comparison between bare sulfur, porous C-S, and porous C-S/GO confirmed the importance of the GO coating for our sulfur-based cathode materials. The complete coating of GO on porous C-S prevented the dissolution of sulfur as polysulphide, and the GO coating also worked as a stable interface between the binder and the active sulfur material. 4.4 Conclusion We have reported a method of fabricating porous C-S/GO composite for cathode material in Li-S batteries, and achieved excellent rate capability and cycling stability, namely 600 mAh/g after 600 cycles at 1 C (1C=1672 mAh/g). And there was only 0.04% capacity decay per cycle between the 2 nd cycle and the 600 th cycle. Our study showed that conductive porous carbon matrix can increase electric conductivity of the electrode and sulfur utilization. Further wrapping with GO can mitigate polysulphide dissolution problem, and provide a chemically stable interface between active material and binder to keep active materials attached to electrodes during cycling. Both conductive porous carbon framework and GO wrapping are indispensable to achieve high-rate, and long- cycling sulfur cathodes in Li-S batteries. 71 4.5 Reference [1] Bruce, P. G.;Freunberger, S. A.;Hardwick, L. J.; Tarascon, J.-M. Li-O-2 and Li-S batteries with high energy storage. Nature Materials 2012, 11, 19-29. [2] Song, M.-K.;Zhang, Y.; Cairns, E. J. A Long-Life, High-Rate Lithium/Sulfur Cell: A Multifaceted Approach to Enhancing Cell Performance. Nano Letters 2013, 13, 5891-5899. [3] Ji, X.; Nazar, L. F. Advances in Li-S batteries. Journal of Materials Chemistry 2010, 20, 9821-9826. [4] Manthiram, A.;Fu, Y.; Su, Y.-S. Challenges and Prospects of Lithium-Sulfur Batteries. Accounts of Chemical Research 2013, 46, 1125-1134. [5] Yang, Y.;Zheng, G.; Cui, Y. 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Nature Nanotechnology 2008, 3, 101-105. [21] Xiao, L.;Cao, Y.;Xiao, J.;Schwenzer, B.;Engelhard, M. H.;Saraf, L. V.;Nie, Z.;Exarhos, G. J.; Liu, J. A Soft Approach to Encapsulate Sulfur: Polyaniline Nanotubes for Lithium-Sulfur Batteries with Long Cycle Life. Advanced Materials 2012, 24, 1176-1181. [22] Zhou, G.;Wang, D.-W.;Li, F.;Hou, P.-X.;Yin, L.;Liu, C.;Lu, G. Q.;Gentle, I. R.; Cheng, H.-M. A flexible nanostructured sulphur-carbon nanotube cathode with high rate performance for Li-S batteries. Energy & Environmental Science 2012, 5, 8901-8906. [23] Guo, J.;Xu, Y.; Wang, C. Sulfur-Impregnated Disordered Carbon Nanotubes Cathode for Lithium–Sulfur Batteries. Nano Letters 2011, 11, 4288-4294. [24] Jayaprakash, N.;Shen, J.;Moganty, S. S.;Corona, A.; Archer, L. A. Porous Hollow Carbon@Sulfur Composites for High-Power Lithium–Sulfur Batteries. Angewandte Chemie International Edition 2011, 50, 5904-5908. [25] Fu, Y.; Manthiram, A. Orthorhombic Bipyramidal Sulfur Coated with Polypyrrole Nanolayers As a Cathode Material for Lithium–Sulfur Batteries. The Journal of Physical Chemistry C 2012, 116, 8910-8915. 73 [26] Fu, Y.;Su, Y.-S.; Manthiram, A. Sulfur–Carbon Nanocomposite Cathodes Improved by an Amphiphilic Block Copolymer for High-Rate Lithium–Sulfur Batteries. ACS Applied Materials & Interfaces 2012, 4, 6046-6052. 74 Chapter 5. Conclusion and Future Work 5.1 Conclusion This thesis is mainly focused on the progress I made in USC on new silicon-based anode material and sulfur-based cathode materials for rechargeable lithium batteries with improved energy density, power density, and cycle life. Specifically, for silicon-based anode material study, a coaxial Si-ATO-Si nanotube array structure was fabricated and used as lithium-ion battery anode in Chapter 2. The coaxial Si-ATO-Si nanotube array demonstrated high specific capacity and excellent cycling performance. After 100 cycles, the capacity still remained above 1500mAh/g and the capacity decay was less than 0.05% per cycle. This excellent cycling stability was due to the unique coaxial structure in which ATO provided a strong inert scaffold, and a rough surface for Si adhesion. This novel structure thus can be a promising candidate for anode material to improve lithium-ion battery performance. For sulfur-based cathode materials study, sulfur particles were coated with GO in a generic method, and tested as Li-S battery cathode in Chapter 3. We found that the GO/c- S composite material exhibits significant improvements in electrochemical performance over c-S particles without coating. Galvanic charge-discharge test using GO/c-S showed 75 that 800 mAh/g specific capacity is retained after 1000 cycles, which is over 6 times larger than that of commercial metal oxide cathode materials (e.g. LiCoO2=120 mAh/g). To further improve the power density of sulfur-based cathode material, we reported a method of fabricating porous C-S/GO composite for cathode material in Li-S batteries, and achieved excellent rate capability and cycling stability, namely 600 mAh/g after 600 cycles at 1 C (1C=1672 mAh/g) in Chapter 4. And there was only 0.04% capacity decay per cycle between the 2 nd cycle and the 600 th cycle. Our study showed that conductive porous carbon matrix can increase electric conductivity of the electrode and sulfur utilization. Further wrapping with GO can mitigate polysulphide dissolution problem, and provide a chemically stable interface between active material and binder to keep active materials attached to electrodes during cycling. Both conductive porous carbon framework and GO wrapping are indispensable to achieve high-rate, and long-cycling sulfur cathodes in Li-S batteries. 5.2 Future Work Although the results presented here have demonstrated the potential of nano-structured silicon-based anode materials and sulfur-based cathode materials in next-generation lithium rechargeable batteries with improved energy density, power density, and cycle life, it could be further developed in at least two ways. First, a full battery using both 76 silicon anode and sulfur cathode has not yet been demonstrated in this thesis. Coupling silicon anode and sulfur cathode together could bring new issues, such as the selection of electrolyte considering different electrolytes were used in silicon and sulfur half-cell studies. Second, lithium needs to be added to either silicon anode or sulfur cathode to make a full battery. 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Abstract (if available)

Abstract Silicon and sulfur are both attractive electrode materials for next-generation rechargeable lithium batteries because of their abundance, high specific capacity, and low cost. This thesis is mainly focused on the research progress I made on silicon-based anode materials and sulfur-based cathode materials for rechargeable lithium batteries with improved energy density, power density, and cycle life. ❧ In Chapter 1, the working mechanism of lithium-ion battery, and commercial anode/cathode materials are introduced, which serve as the background information for the following chapters. Then in Chapter 2, the research progress I made in silicon-based anode is presented. To tackle with the problems of silicon anode, such as poor cyclability and early capacity fading due to significant volume change during lithiation and delithiation process, I fabricated a coaxial silicon / anodic titanium oxide / silicon (Si-ATO-Si) nanotube array structure grown on titanium substrate demonstrating excellent electrochemical cyclability. The ATO nanotube scaffold used for Si deposition has many desired features, such as rough surface for enhanced Si adhesion, and direct contact with the Ti substrate working as current collector. More importantly, the ATO scaffold provides a rather unique advantage that Si can be loaded on both the inner and outer surfaces, and an inner pore can be maintained to provide room for Si volume expansion. This coaxial structure shows a capacity above 1500mAh/g after 100 cycles, with less than 0.05% decay per cycle. ❧ In Chapter 3 and Chapter 4, two kinds of sulfur-based cathode materials are reported. In Chapter 3, a generic and facile method of coating graphene oxide (GO) on sulfur particles is presented. The applications of sulfur/GO core-shell particles as Li-S battery cathode materials are further investigated and the results show that sulfur/GO exhibit significant improvements over bare sulfur particles without coating. Galvanic charge-discharge test using GO/sulfur particles shows a specific capacity of 800 mAh/g is retained after 1000 cycles at 1 A/g current rate if only the mass of sulfur is taken into calculation, and 400 mAh/g if the total mass of sulfur/GO is considered. Most importantly, the capacity decay over 1000 cycles is less than 0.02% per cycle. The coating method developed in this study is facile, robust, and versatile, and is expected to have wide range of applications in improving the properties of particle materials. ❧ In Chapter 4, in order to further improve the power density of sulfur-based cathode material, a method of fabricating graphene oxide (GO) wrapped porous carbon/sulfur composite (porous C-S) for high performance lithium-sulfur (Li-S) battery cathode material is reported. A porous C-S composite using conductive porous carbon as framework and sulfur within its channels as filler is synthesized to generate essential electrical contact to the insulating sulfur, thus achieving high specific capacity. Further graphene oxide wrapping over porous C-S is used to mitigate the problem of intermediate polysulphide dissolution into electrolyte during lithiation/delithiation. The electrochemical performance of GO wrapped porous C-S (porous C-S/GO) is investigated, and the results show that porous C-S/GO exhibits significant improvement over bare sulfur and porous C-S without GO wrapping in terms of specific capacity and cycling stability, respectively. Galvanic charge-discharge test using porous C-S/GO shows that a specific capacity of 600 mAh/g is retained after 600 cycles at 1 C (=1673 mA/g) current rate if the total mass of porous C-S/GO is considered. Our study proves that graphene oxide can provide a chemically stable interface between active material and binder to keep active materials attached to electrodes during cycling. This could serve as an important guideline for future sulfur-based cathode materials design.

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Asset Metadata

Creator Rong, Jiepeng (author)

Core Title Nanostructured silicon anode and sulfur cathode for lithium rechargeable batteries

School Viterbi School of Engineering

Degree Doctor of Philosophy

Degree Program Materials Science

Publication Date 01/29/2015

Defense Date 12/01/2014

Publisher University of Southern California (original), University of Southern California. Libraries (digital)

Tag lithium rechargeable batteries,OAI-PMH Harvest,silicon anode,sulfur cathode

Format application/pdf (imt)

Language English

Contributor Electronically uploaded by the author (provenance)

Advisor Zhou, Chongwu (committee chair), Goo, Edward K. (committee member), Wu, Wei (committee member)

Creator Email jrong@usc.edu

Permanent Link (DOI) https://doi.org/10.25549/usctheses-c3-526687

Unique identifier UC11298601

Identifier etd-RongJiepen-3150.pdf (filename),usctheses-c3-526687 (legacy record id)

Legacy Identifier etd-RongJiepen-3150.pdf

Dmrecord 526687

Document Type Dissertation

Format application/pdf (imt)

Rights Rong, Jiepeng

Type texts

Source University of Southern California (contributing entity), University of Southern California Dissertations and Theses (collection)

Access Conditions The author retains rights to his/her dissertation, thesis or other graduate work according to U.S. copyright law. Electronic access is being provided by the USC Libraries in agreement with the a...

Repository Name University of Southern California Digital Library

Repository Location USC Digital Library, University of Southern California, University Park Campus MC 2810, 3434 South Grand Avenue, 2nd Floor, Los Angeles, California 90089-2810, USA